Thin film solid-state batteries are widely used in medical devices to provide power for microelectromechanical systems (MEMS), and as on-chip power sources because of their light weight, small size, and high power storage density [Ref. 67-69]. Rechargeable thin-film Li/Li-ion batteries have been widely investigated as higher energy density replacements for nickel-metal hydride rechargeable batteries [Ref. 1] and also for use as on-chip power sources [Ref. 2]. Secondary, or rechargeable, thin-film Li/Li-ion batteries are multi-layer structures [Ref. 2, 3]. These batteries consist of a current collector, a cathode, an anode and a solid electrolyte. A schematic illustration of such a device based upon the storage of lithium ions is shown in FIG. 1. Their total thickness is normally less than 15.0 μm including the ˜6 μm thick environmental protective layer that is needed to avoid absorption of oxygen and moisture [Ref. 4]. During the discharge of these batteries, lithium ions diffuse from the anode, migrate through the solid electrolyte and accumulate in the cathode layer by intercalation [Ref. 5]. Electron flow occurs in the reverse direction through a load resistor. Both processes reverse during a charging cycle.
The effective charge and discharge of rechargeable thin-film Li/Li-ion batteries requires a thin-film electrolyte that blocks electron transport while allowing high conductive passage of lithium ions [Ref. 6]. This translates into an electron conductivity less than ˜10−14 S/cm and a Li-ion conductivity in the 10−5-10−8 S/cm range [Ref. 5]. It is also essential that the thin-film electrolyte through thickness be free of interconnected porosity or cracks to avoid electrical breakdown (electrical shorting) when a voltage is applied across the electrolyte [Ref. 7]. The thin-film electrolyte also needs to be both thin (1-2 μm) to reduce internal resistive losses, and uniform in thickness to avoid localization of the discharge process [Ref. 6]. In rechargeable thin-film lithium batteries, the thin-film electrolyte should also not decompose when in intimate contact with a lithium metal anode. The thin-film electrolyte therefore plays a crucial role in the operation of rechargeable thin-film Li/Li-ion batteries and its synthesis must be carefully controlled to achieve the desired composition and structure needed to optimize battery performance [Ref. 6].
Many electrolyte chemistries have been investigated for rechargeable thin-film Li-ion batteries [Ref. 8, 9]. They include Li2S—P2S5—LiI [Ref. 10], Li2S—SiS2—LiI [Ref. 11], Li2S—SiS2—Li3PO4 [Ref. 12], Li2O—P2O5—Li2SO4 [Ref. 13], Li2O—B2O3—LiI [Ref. 14], Li2O—Al2O3—B2O3 [Ref. 15], Li2O—Al2O3—SiO2 [Ref. 16], Li2O—SiO2—B2O3 [Ref. 17], and lithium phosphorous oxynitride (Lipon) [Ref. 18]. While the Li-sulfate electrolytes have a high Li-ion conductivity of 10−3-10−4 S/cm, they are highly reactive with air and are difficult to fabricate because of their corrosive nature. The Li-oxide electrolytes are much more stable in oxygen but also absorb moisture [Ref. 9]. They also have a lower Li-ion conductivity (10−6-10−8 S/cm) than Li-sulfate electrolytes [Ref. 8, 9]. Furthermore, both the Li-sulfate and Li-oxide electrolytes are decomposed when they come into contact with a lithium anode with an applied potential of up to 5.0 V [Ref. 9]. Lipon electrolytes have been extensively investigated for thin film Li/Li-ion battery applications [Ref. 3], and are widely used in thin film Li/Li-ion batteries because they do not decompose when in contact with a lithium anode. They also possess a relatively high Li-ion conductivity (in the 10−6-10−7 S/cm range) [Ref. 18, 19]. Since the ionic conductivity of amorphous Lipon films is generally more isotropic and higher than that of crystalline films, amorphous Lipon films are preferred for solid electrolyte applications [Ref. 6].
Rechargeable thin-film batteries based upon lithium anodes and Lipon electrolytes have been fabricated by a combination of resistive thermal evaporation and reactive RF-magnetron sputtering. Bate et al. successfully synthesized Lipon films using RF-magnetron sputtering under either mixed Ar—N2 or pure N2 atmospheres [Ref. 18]. While high quality Lipon films could be synthesized, the RF-magnetron sputtering suffered from a very low deposition rate (˜1 nm/min.) due to the low working pressure (˜20 mTorr) and low power (12-40 W) needed to avoid cracking of the target [Ref. 20]. Attempts have been made to increase the deposition rate of Lipon films by using a N2—He (instead of an Ar) plasma. However, the deposition rate of the Lipon films still remained less than 3.0 nm/min. [Ref. 21]. Several attempts have been made to synthesize Lipon films by other deposition approaches including Pulse Laser Deposition (PLD) [Ref. 22], Ion Beam Assisted Deposition (IBAD) [Ref. 23], and electron-beam (EB) evaporation [Ref. 24]. The deposition rate of PLD films was in the 13.3-50 nm/min range while those synthesized using IBAD could be grown at up to ˜66 nm/min. Both approaches therefore enabled film growth at much higher deposition rates than the RF-magnetron sputtering approach. The PLD and IBAD methods resulted in films with a Li-ion conductivity in the 1.4×10−6-4×10−8 S/cm range [Ref. 22, 25]. Unfortunately, the Lipon films deposited by the PLD approach had a very rough surface morphology while those synthesized by the IBAD approach contained large tensile stresses which led to film cracking and electrical shorting in metal/Lipon/metal test cells [Ref. 22, 23].
The EB evaporation approach has employed moderate power (300 W) e-beams for the evaporation of a Lipon source and a moderate power (˜250 W), 13.54 MHz inductively coupled Ar—N2 plasma (ICP) reactor for reactive synthesis in an ionized nitrogen environment [Ref. 24]. These EB evaporated Lipon films had a Li-ion conductivity of ˜10−7-10−8 S/cm and could be grown at somewhat higher deposition rates (˜8.33 nm/min) than those achievable using reactive RF-magnetron sputtering [Ref. 24]. While the maximum deposition rate for this approach was significantly less than that of the PLD and IBAD approaches, it appears a promising route for the more economical deposition of Lipon films, especially if the deposition rate can be improved without adversely affecting other properties of the electrolyte.
After depositing the cathode and Lipon films by RF-magnetron sputtering, a battery structure is usually completed by deposition of a lithium anode using a resistive thermal evaporation method [Ref. 3]. Lithium alloys are highly reactive with water vapor and so the resistive thermal evaporation step is usually conducted inside a dry, inert environment in order to avoid reactions during the transfer of samples [Ref. 26].
A plasma-assisted directed vapor deposition (PA-DVD) approach for the deposition of various metals and metal oxides has recently been developed [Ref. 27, 28]. This electron beam evaporation-based approach has a multi-source capability raising the possibility of depositing all the layers of a thin film battery within a single reactor [Ref. 27]. The deposition techniques also allow uniform conformal coating of surfaces including the interior of cellular structures.
During discharge, Li+ ions leave the anode, diffuse through the electrolyte, and are intercalated within the cathode material. An electron current simultaneously flows via the leads between the cathode and anode through a load resistor. The overall discharge reaction can be written as:Anode: xLi→xLi++xe−Cathode: M+xLi++xe−→LixMwhere M is a transition metal atom in a transition metal oxide that typically serves as the host cathode for the lithium ions.
Lithium transition metal oxides such as LiCoO2 and LiMn2O4 are widely used for the cathode material [Ref. 70-72]. LiCoO2 has a layered structure (R-3m), which facilitates lithium insertion and extraction during battery operation [Ref. 70]. It is widely used in commercial batteries in part because of its high specific charge storage capacity (˜130 Ah/kg) and excellent rechargeability (>1000 cycles) [Ref. 70-72]. However, LiCoO2 is costly, and it has significant toxicity issues [Ref. 73]. LiMn2O4 is a candidate alternative cathode material for high energy density battery applications [Ref. 74, 75].
The performance of the cathode layers in thin films batteries depends on many aspects of the film including its composition [Ref. 76-80], degree of crystallinity [Ref. 81], grain size [Ref. 82], and the film's pore volume fraction and topology [Ref. 83, 84]. Stoichiometric LiMn2O4 has a cubic spinel structure (Fd3m), FIG. 19 [Ref. 85]. FIG. 19 shows the unit cell of cubic spinel LiMn2O4. The space group of the cubic spinel is Fd3m, and there are 56 atoms per unit cell. Lithium atoms are tetrahedrally coordinated with oxygen atoms, while the manganese atoms are octahedrally coordinated with oxygen atoms. In the ideal cubic spinel structure, the Li and Mn ions are located at tetrahedral (8a) and octahedral (16d) sites in the cubic-closed-packed oxygen ion frame (32e) [Ref. 86]. Li+ ions can be intercalated and deintercalated reversibly in this lattice during charging and discharging without breaking the basic lattice structure [Ref. 86].
The spinel phase of lithium manganese oxide exists over a wide composition range [Ref. 87] and can be relatively easily grown if the appropriate synthesis conditions are used [Ref. 86]. By changing the growth temperature and oxygen partial pressure, both stoichiometric Li1+xMn2−xO4 [Ref. 76, 78], and non-stoichiometric LiMn2O4±x [Ref. 79, 88] films have been grown. The value of x in these lithium manganese oxide films modifies electrochemical properties, and changes both the specific storage capacity and the cylclic rechargeability [Ref. 76].
These changes in properties are thought to result from modifications to the average Mn oxidation state required to achieve charge neutrality of an overall unit cell [Ref. 76]. Higher Mn oxidation states appear to promote retention of charge storage during repeated cycling, but they reduce the maximum charge storage capacity [Ref. 76]. Better cathode performance is exhibited by films that contain a significant pore volume fraction [Ref. 83, 84]. A small grain size also appears favorable for enhancing the ionic mobility [Ref. 89]. Finally, films with a strong (111) preferred orientation provide improved electrode performance because the (111) direction is a channeling direction for Li ions, and enhance their effective diffusivity is enhanced in this direction [Ref. 83, 90].
Lithium-transition metal oxide films such as lithium manganese oxide can be fabricated by either electron-beam evaporation [Ref. 82, 89, 91], sputtering [Ref. 81, 90-94], or by pulsed laser deposition [Ref. 80, 95-97]. For example, reactive electron beam evaporation has been used to synthesize LiMn2O4 films with a small grain size and good electrochemical performance [Ref. 89]. Wang et al have used magnetron sputtering to grow Li4Ti5O12 thin films, which are an isostructure to LiMn2O4 [Ref. 83]. They have produced two types of surface morphologies of films; one relatively dense and the other consisting of island-like grains with interconnected grain boundary pores. In addition, they examined the electrochemical properties of different degree of textures with the same surface morphologies. They found that the highly textured, porous films with interconnected grain boundary pores exhibited better electrode performance. The typical deposition rate of sputtering techniques is usually less than ˜10 nm/min [Ref. 93].
An electron-beam directed vapor deposition (EB-DVD) technique has recently been developed for synthesizing binary metal oxides with controlled compositions and pore fraction [Ref. 98-102]. In this approach, an electron beam is used to evaporate a source material located in a water cooled crucible positioned in the throat of a nozzle that forms a supersonic gas jet. This jet entrains and transports the vapor to a substrate. The gas jet speed is determined by the pressure difference between the pressure upstream of the nozzle opening and that downstream in the growth chamber, and by the ratio of the specific heats of the gas used to form the jet [Ref. 103]. Inert carrier gases such as helium or argon sometimes doped with small amount of oxygen or other reactive gases are typically used. Changing the upstream to downstream modifies the degree of collimation of the vapor flux and the fraction of vapor deposited on a substrate [Ref. 99]. The technique has been used to grow thick yttrium stabilized zirconia coatings at very high deposition rates in excess of ˜10 μm/min for use as thermal barriers [Ref. 98, 99, 104]. The invention described shows that the directed vapor deposition approach can be used to deposit the cathode layers of thin film batteries with quite compositionally complex chemistries. Using the lithium manganese oxide system as an example we show that the film composition, phase content, texture and pore volume fraction can all be controlled by the jet speed and the deposition pressure [Ref. 98, 99].